Armour steel

ABSTRACT

A low-carbon martensitic armour steel comprises at least Fe, C, Si and Ni and has a ratio of yield strength to ultimate tensile strength of less than 0.7. The steel includes retained austenite at a volume fraction of at least 1%. The low-carbon martensitic armour steel can be prepared by subjecting a steel which comprises at least Fe, C, Si and Ni and which has a martensite start temperature of less than 210° C. to an austenisation heat treatment step at a temperature of at least 800° C., quenching the steel, and subjecting the steel to a tempering step at a temperature of less than 300° C.

THIS INVENTION relates to armour steel. In particular, the invention relates to a low-carbon martensitic armour steel, and to a method of preparing a low-carbon martensitic armour steel.

Over the last few decades, military and security specifications have been developed for armour steel plates based on their predicted behaviour when impacted by high velocity rounds. Hitherto the hardness and the tensile and impact properties (in particular, the yield strength, the ultimate tensile strength, the elongation at room temperature and the transverse Charpy-V impact energy at −40° C.) were the main design parameters for most armour steel plates. Current specifications for military and security applications recommend a minimum Brinell hardness range of 540 to 600 BHN or 55 to 60 Rockwell C, and a minimum yield and ultimate tensile strength of 1500 MPa and 1700 MPa respectively, while a minimum Charpy-V impact energy of 13 Joules at −40° C. and a minimum elongation of 6% (50 mm gauge length) is required in South Africa. These requirements have been successfully achieved over many years by a suitable combination of chemical composition and heat treatment parameters of armour steels for plates generally thicker than 8.5 mm up to 30 mm.

The inventors believe that neither a higher hardness nor higher mechanical properties are exclusive or even reliable criteria for predicting the ballistic performance of martensitic armour steels. Instead, the ratio between the yield and the ultimate tensile strength, rather than the hardness, is believed by the inventors to be more important in approximating the true dynamic fracture and spalling strength of the armour steel.

According to one aspect of the invention, there is provided a low-carbon martensitic armour steel comprising at least Fe, C, Si and Ni which has a ratio of yield strength to ultimate tensile strength of less than 0.7, and which includes retained austenite at a volume fraction of at least 1%.

Preferably, the ratio of yield strength to ultimate tensile strength is less than or equal to 0.65, more preferably less than or equal to 0.60.

The armour steel may comprise 0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni, with the balance being mostly Fe. Preferably, the Si content is higher than 0.8 weight %.

The armour steel may comprise also Mn, Cr and Mo. Thus, the armour steel may comprise 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight % Mo. Instead of, or in addition to, Mo, the armour steel may comprise Co and/or W.

The armour steel may have a martensite start temperature of less than 210° C., preferably less than 200° C.

The armour steel may have a plate thickness of from 4.5 mm to 8 mm, typically between 4.5 mm and 6 mm.

Typically, the armour steel has a volume fraction of retained austenite of less than 7%.

The armour steel may have a micro-structure in which the martensite is predominantly present as twinned plate martensite and not lath martensite. Preferably, the twinned plate martensite is combined with retained austenite in the same micro-structure.

The armour steel typically includes cementite. Any cementite which is present is preferably predominantly present as dispersed particles and not as coarse strings.

The armour steel is typically in the form of a plate with a thickness δmm. The armour steel may have a Ballistic Parameter BP of at least 0.01 where BP=volume fraction of retained austenite/exp(δ).

The armour steel may have a Charpy-V impact energy of less than 13 Joules at −40° C. The armour steel may have a hardness which is less than 600 BHN or less than 640 VHN. The armour steel may have a yield strength of less than 1700 MPa, e.g. less than or equal to 1500 MPa. The armour steel may have an ultimate tensile strength of less than 2000 MPa, e.g. less than 1700 MPa.

According to another aspect of the invention, there is provided a method of preparing a low-carbon martensitic armour steel, the method including

subjecting a steel which comprises at least Fe, C, Si and Ni and which has a martensite start temperature of less than 210° C. to an austenisation heat treatment step at a temperature of at least 800° C.;

quenching the steel; and

subjecting the steel to a tempering step at a temperature of less than 300° C.

The austenisation heat treatment step may be at a temperature of between 870° C. and 950° C. The steel may be subjected to the austenisation heat treatment step for a period of between 20 minutes and 60 minutes.

The tempering step may be at a temperature of between 150° C. and 250° C., e.g. between 150° C. and 200° C.

The steel may be subjected to the tempering step for a period of between 20 minutes and 60 minutes, e.g. between 20 minutes and 30 minutes.

Typically, the steel is in the form of a plate with a thickness of from 4.5 mm to 8 mm.

The steel may comprise 0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni, the balance being mostly Fe. Preferably, the steel has a Si content of higher than 0.8 weight %.

The steel may comprise also Mn, Cr and Mo. Thus, the steel may comprise 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight % Mo. Instead of, or in addition to, Mo, the steel may comprise Co and/or W.

The steel may have a martensite start temperature of less than 200° C.

The invention extends to a low-carbon martensitic armour steel produced by the method as hereinbefore described.

The invention will now be described in more detail, by way of the following experimental studies and the drawings, and with reference to five experimental steels G1A, G1B, G2A, G2B and G3, and three conventional steels A66, M38 and RL5.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows contours of constant YS/UTS ratio for the steel G1A;

FIG. 2 shows contours of constant tensile strength (in MPa) of the steel G1A;

FIG. 3 shows contours of constant Charpy impact energy (CIE in Joules) at −40° C. for the steel G1A;

FIG. 4 shows lines of constant Vickers hardness of the steel G1A;

FIG. 5 shows three-D surfaces representing the Charpy impact energies of the experimental steels G1A, G1B, G2A and G3;

FIG. 6 shows three-D surfaces representing the ultimate tensile strengths of the steel G1A that passed ballistic testing, and the steel G2B that failed the ballistic testing;

FIG. 7 shows three-D representations of the ratios YS/UTS of the experimental steels showing lower values for the steels G1A and G1B that passed the ballistic testing, and higher values for the steels G2B and G3 that failed the ballistic testing;

FIG. 8 shows secondary electron scanning microscopy of the shear lips of steel G1B tempered at 300° C., showing the cavities and decohesion around MnS particles upon Charpy impact testing at −40° C.;

FIG. 9 shows secondary electron SEM of fracture surface showing an elongated plate-like inclusion of manganese sulphide, observed after the tensile test of steel G1B tempered at 150° C., with large cavities around the inclusions of the MnS;

FIG. 10 shows secondary electron SEM of fractured surface of untempered steel G2A after ballistic testing showing no decohesion at the interface between matrix and MnS particle;

FIG. 11 shows secondary electron SEM of fractured surface of untempered steel G2A after ballistic testing showing striations that suggest a propagation of the fracture front under variable tensile stress upon ballistic impact;

FIG. 12 shows thin foil transmission electron microscopy of the steel G1A after tempering at 180° C. and before ballistic testing, with a scale bar length of 500 nm;

FIG. 13 shows thin foil transmission electron microscopy of the steel G2A after tempering at 180° C. and before ballistic testing, with a scale bar length of 500 nm;

FIG. 14 shows thin foil transmission electron microscopy of the steel G2B after tempering at 180° C. and before ballistic testing, with a scale bar length of 500 nm;

FIG. 15 shows a bright field thin foil transmission electron micrograph of the steel G1A after tempering at 300° C.;

FIG. 16 shows a bright field thin foil transmission electron micrograph of the steel G1B after tempering at 300° C.;

FIG. 17 shows a bright field thin foil transmission electron micrograph of the steel G1A (0.21% Si) tempered at 400° C. showing large strings of cementite; and

FIG. 18 shows a bright field thin foil transmission electron micrograph of the steel G1B (1.06% Si) tempered at 400° C. showing coarse cementite.

STUDY 1

In a series of experimental tempered martensitic steel alloys it was observed that for a given chemical composition, the heat treatment parameters for advanced ballistic performance are different from those required for higher mechanical properties, rendering the often specified relationship between mechanical properties and ballistic performance questionable. Systematic analysis of the microstructures and the fracture surfaces of thirteen laboratory melted tempered martensitic armour plate steels were carried out to understand the improved ballistic performance of these steels of which the mechanical properties were actually lower than currently specified for military and security applications. It was, furthermore, observed that the detrimental effect of inclusions on ballistic performance depends on the tempering temperature and on the strain rate.

In this study a comparison of the effect of manganese sulphide inclusions in “slower” strain rate Charpy impact tests at −40° C. and tensile tests at room temperature, and in “higher” strain rate ballistic tests on the fracture mode of plates was made to explain the discrepancies between the performance predictions based on mechanical properties as in many current design specifications, and the observed ballistic performance for plates of tempered martensitic steels with thicknesses less than 8.5 mm.

Localised thin foil transmission electron microscopy (TEM) of ballistic impacted regions of these steels suggested the presence of high temperatures during the impact that induced phase transformations from an initial twinned martensite to austenite and back to an untwined martensite.

Materials and Experiments

Chemical Composition and Manufacturing

Initially five experimental armour steels, namely steels G1A through to G3 were subjected to standard ballistic testing and their performance compared to those of three currently produced and used armour steels, here named A66, M38 and RL5. Their chemical compositions are shown in Table 1.

TABLE 1 Chemical compositions in wt % of the initial five experimental martensitic armour steels and three current commercially produced armour steels Steel C Mn P S Si Cu Ni Cr Mo V Nb Ti N G1A 0.39 1.22 0.008 0.003 0.21 0.10 2.99 1.49 0.5 0.006 0.002 0.003 0.0049 G2A 0.37 1.15 0.015 0.011 1.06 0.14 3.8 0.52 0.43 0.008 0.008 0.007 0.0036 G2B 0.39 0.65 0.017 0.009 0.8 0.23 2.8 0.22 0.24 0.003 0.006 0.01 0.0051 G2A 0.37 0.40 0.016 0.011 0.43 0.33 2.3 0.24 0.3 0.006 0.006 0.009 G3 0.34 0.39 0.019 0.012 0.40 0.32 2.43 0.27 0.37 0.009 0.009 0.008 A66 0.37 0.68 0.003 0.002 0.24 0.005 1.9 0.48 0.32 0.004 0.001 0.003 0.009 M38 0.38 0.55 0.004 0.002 0.77 0.1 1.79 0.14 0.36 0.001 Nil 0.007 RL5 0.32 0.86 0.008 0.002 0.18 0.26 2.8 0.79 0.45 0.005 0.001 0.003 0.009

The 5 kg vacuum melted alloys were cast into a 45 mm×70 mm×230 mm mild steel mould, the ingots were solution treated for one hour at 1100° C. before hot rolling in four passes of 20% reduction each down to 6±0.2 mm thickness at a finishing temperature of between 950 and 900° C. and then air cooled. The plates were then austenitised at 900° C. for 20 minutes, water quenched and tempered at 180° C. for one hour. The microstructures and the phases present before ballistic testing were analysed by thin foil transmission electron microscopy and X-ray diffraction. The plate's sizes for ballistic testing were 200 mm to 250 mm wide and 500 to 550 mm in length.

Ballistic Testing

The plates of these steels were ballistically tested according to the NATO and to the South African specifications. The plates were clamped on one of the 250 mm edges and passed the ballistic test if they withstood 5.56 mm rounds fired at a measured striking velocity higher than 940±10 m/s, from a distance of 30 metres and 0° obliquity angle in a ballistic tunnel. The parameters and effect of each impact were recorded and in general, five rounds were fired at each plate. No light should be visible through any of the impacted regions on a plate for a successful test.

Results of the Ballistic Testing

Ballistic Report

Table 2 shows the measured mechanical properties of the steels used in characterising the armour materials according to the current specifications. In addition, the martensite start temperatures (M_(s)), the volume fractions of retained austenite (% RA) and the yield to tensile strength ratios (YS/UTS) of the plates, were also considered as design parameters.

TABLE 2 Properties of the plates before ballistic testing Impact Retained energy Thickness Hardness YS UTS austenite Elongation at −40° C. Ms [° C.] [mm] VHN (30 kg) [MPa] [MPa] YS/UTS [%] [%] [Joules] measured G1A 6.0 ± 0.2 578 880 1650 0.53 6 4 10 196 G1B 6.0 ± 0.2 565 960 1700 0.56 4 6 13 210 G2A 6.0 ± 0.2 610 1500 2200 0.68 0.6 8 14 255 G2B 6.0 ± 0.2 520 1500 2000 0.75 0.6 12 17 271 G3 6.0 ± 0.2 490 1300 1700 0.76 0.5 14 18 309 A66 6.5 ± 0.2 640 1300 1900 0.68 0.6 6 16 (full 243 size) M38 8.5 ± 0.2 620 1300 1800 0.72 0.6 7 16 (full 243 size) RL5  12 ± 0.2 540 1400 1700 0.82 <0.5 6 18 (full 285 size)

The Charpy impact energy was measured at −40° C. on sub-sized specimens of 5×10×55 mm due to the thickness limitation of the heat treated plates. The Charpy impact energies of the armour steels A66, M38 and RL5 are those specified for full size specimens by the respective manufacturers.

The results on the ballistic performance of these eight martensitic steels are presented in Table 3.

TABLE 3 Results of the ballistic testing Firing Steel Distance Ballistic designation [m] performance G1A (6.0 mm) 30 Passed G1A (6.0 mm) 10 Passed G1B (6.0 mm) 30 Passed G2A* (6.0 mm) 30 3 out of 5 failed G2B** (6.0 mm) 30 Failed G3** (6.0 mm) 30 Failed A66 (6.5 mm) 30 Passed M38 (8.5 mm) 30 Passed RL5** (12 mm) 30 Failed

In the rest of this patent specification, the steels will be designated with (i) no asterisk=passed (ii) one asterisk=partially failed and (iii) two asterisks=failed the ballistic test.

Industrial experience shows that plates of steels A66 and M38 are not reliable for thicknesses smaller than indicated in Table 3.

From the ballistic report it appears that thinner martensitic plates having lower values of the ratio YS/UTS are more efficient in resisting ballistic perforation. It, however, also appears that the higher hardness, higher tensile strength and ultimate tensile strength are not necessarily appropriate in indicating good ballistic performance for these martensitic steels.

Measured Mechanical Properties

Sample Preparation and Measurement of the Mechanical Properties

The tensile and the sub-sized Charpy specimens of the steels were austenitised for 30 minutes at 800° C., 850° C., 900° C. and 950° C. respectively in an Argon atmosphere and were then water quenched to form the martensitic microstructure. Tempering treatments of the specimens were carried out at 150, 180, 200, 250, 300, 350 and 400° C. for different times between 15 and 60 minutes.

The tensile specimens were wire cut parallel to the rolling direction and the yield strength, the ultimate tensile strength and the elongation were determined using an INSTRON 8500 hydraulic tensile testing machine. The ultimate tensile strength at room temperature is considered as the first constraint on the ballistic performance according to some current design procedures.

As in the case of the tensile properties, current specifications exist for the Charpy impact energy of armoured plate steels, a test that is at a relatively “slower” strain rate than that during ballistic testing. The Charpy impact energy of the sub-sized specimens was measured at −40° C. to construct the second constraint on the ballistic performance.

The dependence of the mechanical properties and ballistic performance of these armour plate steels on the M_(s) temperature was also examined, as measured in a THETA 734 Single Silica Push Rod LVDT dilatometer. The thin foils for Transmission Electron Microscopy were prepared from the 3 mm diameter discs electro-tube-cut from the armour plates in the quenched and in the quenched and tempered conditions before and after ballistic impact.

The fracture surfaces after tensile, Charpy impact testing, and after ballistic testing were analysed in secondary electron mode on a JEOL JSM-6300 scanning electron microscope to determine the mode of cleavage and the possible role of inclusions in the fracture mechanism.

Ultimate tensile strength, yield strength to ultimate tensile strength ratio and Charpy Impact Energy

The ultimate tensile strength (UTS), the yield strength to ultimate tensile strength ratio YS/UTS and the Charpy impact energy at −40° C. (CIE) of the five steels G1A through to G3 were measured for different austenitisation and tempering temperatures. The equation describing the variations of the these three properties with respect to the heat treatment temperatures were determined for each of the five steels by surface fitting to the measured values.

It was observed that third degree polynomials fitted the results within the experimental ranges of the austenitisation and the tempering temperatures. This led to the general mathematical expression of the curve-fitted surfaces using normalising functions for the two temperatures as defined in equation (1) for the tempering temperature and in equation (2) for the austenitisation temperature. Normalised temperatures were used to minimise the rounding errors in determining the fitting equations.

$\begin{matrix} {T_{tn} = \frac{\left( {T_{t} - T_{tm}} \right)}{\left( {T_{tm} - 25} \right)}} & (1) \\ {T_{an} = \frac{\left( {T_{a} - T_{am}} \right)}{\left( {T_{am} - 800} \right)}} & (2) \end{matrix}$

where T_(t) is the actual tempering temperature in degrees Celsius, T_(tm) is a mean tempering temperature calculated as T_(tm)=(25+300)/2=162.5° C. T_(a) is the actual austenitisation temperature in degrees Celsius and T_(am) is the mean austenitisation temperature calculated as T_(am)=(800+950)/2=875° C.

The correspondence between normalised and actual values of the austenitisation and tempering temperatures is given in Table 4.

TABLE 4 Correspondence between normalised and actual temperatures in degrees Celsius Normalised temperature −1 −0.75 −0.5 −0.25 0 0.25 0.5 0.75 1 Aus- 800 819 837 856 875 894 913 931 950 teniti- sation temper- ature Tem- 25 60 94 128 163 197 231 266 300 pering tem- perature

The particular mechanical property (MP) is then fitted by the equation:

MP(T _(an) ,T _(tn))=a(T _(an))×T _(tn) ³ +b(T _(an))×T _(tn) ² +c(T _(an))×T _(tn) +d  (3)

where the fitting parameters a, b, c and d are polynomials in T_(an) and are of the general form:

p=A×T _(an) ³ +B×T _(an) ² +C×T _(an) +D  (4)

A, B, C and D are constants.

Combining equations (3) and (4) gives a sixth order non-linear equation for the mechanical property in terms of the normalised temperatures T_(an) and T_(tn).

MP(T _(an) ,T _(tn))=(A ₁ ×T _(an) ³ +B ₁ ×T _(an) ² +C ₁ ×T _(an) +D)×T _(tn) ³+(A ₂ ×T _(an) ³ +B ₂ ×T _(an) ² +C ₂ ×T _(an) +D ₂)×T _(tn) ²+(A ₃ ×T _(an) ³ +B ₃ ×T _(an) ² +C ₃ ×T _(an) +D)×T _(tn)+(A ₄ ×T _(an) ³ +B ₄ ×T _(an) ² +C ₄ ×T _(an) +D ₄)  (5)

Mechanical properties of Steel G1A (M_(s)=196° C.)

Fitting Function for the Ultimate Tensile Strength UTS, the YS/UTS and the CIE

For illustration purpose the fitted equations of UTS, YS/UTS and (CIE) at −40° C. of steel G1A are represented mathematically by the functions (6) to (8):

$\begin{matrix} {{UTS} = {{\left( {{94.905T_{sn}^{3}} - {51.94T_{an}^{2}} - {167.71T_{an}} + 137.43} \right) \times T_{tn}^{3}} + {\left( {{{- 177.77}T_{an}^{3}} - {30.043T_{an}^{2}} + {333.18T_{an}} - 302.81} \right) \times T_{tn}^{2}} + {\left( {{{- 204.53}T_{an}^{3}} + {226.44T_{an}^{2}} + {114.32T_{an}} - 203.98} \right) \times T_{tn}} + \left( {{268.71T_{an}^{3}} + {81.169T_{an}^{2}} - {472.81T_{an}} + 1875.9} \right)}} & (6) \\ {\frac{YS}{TS} = {{\left( {{{- 0.0317}T_{an}^{3}} + {0.0012T_{an}^{2}} + {0.0383T_{an}} + 0.0213} \right) \times T_{tn}^{2}} + {\left( {{{- 0.0157}T_{an}^{3}} + {0.0282T_{an}^{2}} + {0.0158T_{an}} + 0.0352} \right) \times T_{tn}} + \left( {{0.0232T_{an}^{3}} + {0.0209T_{an}^{2}} - {0.0584T_{an}} + 0.4607} \right)}} & (7) \\ {{{CIE}\left( {{- 40}{^\circ}\mspace{14mu} {C.}} \right)} = {{\left( {{0.2131T_{an}^{3}} - {0.6036T_{an}^{2}} + {0.5825T_{an}} + 0.0983} \right) \times T_{tn}^{3}} + {\left( {{0.3463T_{an}^{3}} + {0.6073T_{an}^{2}} - {1.0607T_{an}} + 0.5168} \right) \times T_{tn}^{2}} + {\left( {{{- 0.8878}T_{an}^{3}} + {0.6606T_{an}^{2}} - {0.6379T_{an}} + 1.9633} \right) \times T_{an}} + \left( {{{- 0.4262}T_{an}^{3}} - {0.0011T_{an}^{2}} - {0.6425T_{an}} + 6.8781} \right)}} & (8) \end{matrix}$

The objective function may be written in the form:

YS/UTS≦r ₀  (9)

where r₀ is 0.68 for the plates of thicknesses smaller than 6.5 mm, referring to the ballistic report in Table 2.

The optimum regions for the properties represented by these equations are found graphically using two-dimensional projections of contours of equal height (i.e. iso-lines in the normalised (T_(an),T_(tn)) planes) as illustrated in FIGS. 1 to 3 for the steel G1A.

From FIG. 1, it appears that the YS/UTS ratio increases as the tempering temperature is increased. Furthermore, it is inferred from FIG. 2 that a tensile strength higher than 1700 MPa is obtained for steel G1A when the normalised austenitisation temperature is lower than −0.1 (T=867° C.) and the tempering temperature lies between the normalised values of −0.5 and 1, (or actually between 95° C. and 300° C.). If the tempering temperature were lower than 95° C., the tensile strength would become difficult to determine because of the brittle behaviour of this steel in that condition. From FIG. 3, a Charpy impact energy at −40° C. that is higher than the specified 13 Joules, is obtained for the normalised temperatures between −1 and 0.2 or lower than an actual 890° C. and the normalised tempering temperature is above 0.9 or actually above 286° C. The summary of this discussion is presented in Table 5.

TABLE 5 Heat treatment conditions predicted to be favourable for the mechanical properties for steel G1A. Favourable conditions Austenitisation Property temperature Tempering temperature Low YS/TS 840° C. to 950° C. <217° C. High UTS <860° C. 95° C. to 300° C. High CIE(−40° C.) <890° C. >286° C.

The optimum heat treatment region of steel G1A for ballistic application could, therefore, be fixed at austenitisation temperatures higher than 840° C., and lower than 860° C. for the high strength cases. However, it is more difficult to find a compromise between the YS/UTS ratio and the Charpy impact energy at −40° C. in terms of the tempering temperature.

The hardness of steel G1A varies with both the austenitisation temperature and the tempering temperature. The regression analysis and the surface fitting were developed following the same procedure as adopted earlier. It was noted that the hardness of steel G1A decreases very fast to values as low as 450 VHN when the tempering temperature was above 200° C. That would be very low compared to the value of 600 BHN or 640 VHN specified for military and security applications in some countries, e.g. South Africa. The variation of the Vickers hardness with the austenitisation and the tempering temperatures is then written in terms of the normalised temperatures as follows:

VH=(28.123T _(an) ³+46.048T _(an)−26.241T _(an)+31.589)×T _(tn) ³+(−38.351T _(an) ³−33.588T _(an) ²+39.306T _(an)+0.5894)×T _(tn) ²+(−24.576T _(an) ²−1.1265T _(an)−112.31)×T _(tn)+(32.687T _(an) ²+26.939T _(an)+531.32)

The lines of constant Vickers hardness of the steel G1A are shown in FIG. 4.

The hardness of steel G1A in the quenched condition is relatively constant when the austenitisation temperature is increased between 800° C. and 900° C. Above this austenitisation temperature range, for instance at 950° C., however, the maximum hardness that was attained increased. This increase in hardness is mainly due to two factors; firstly the solid solution hardening of the parent austenite due to the increased dissolution of some carbides when the austenitisation temperature was increased and secondly, to a subsequent decrease in the M_(s) temperature leading to a harder untempered martensite with a greater amount of carbon in solution. At 950° C. grain growth of the austenite can also become significant. The decrease in the martensite start temperature may, however, lead to an increase in the volume fraction of retained austenite and imposes a limit on the increase of the hardness of the steel. The grain size of steel G1A after austenitisation for 30 minutes, as determined by the line intercept method using the line scanning function of the scanning electron microscope, increased from 7.0±0.8 μm when the austenitisation temperature was 850° C., to 10±0.8 μm when the austenitisation temperature was 950° C.

The highest tensile strengths are achieved when the austenitisation temperature is below 867° C. and it dropped again above this austenitisation temperature. This effect may also be related to grain growth and the increase in the volume fraction of the retained austenite. Therefore, it appears that both the tensile strength and the hardness increase first with an increase in the austenitisation temperature, but the upper limit in the tensile strength occurs earlier than for the hardness.

The rate of decrease in hardness of steel G1A upon low-temperature tempering, appears to be slower when the austenitisation temperature is lower within the range from 800° C. to 900° C. This trend indicates that at higher austenitisation temperatures the amount of carbon dissolved in the parent austenite is high, which leads to a higher activity of carbon in the martensite upon tempering. The sudden change of slope of the hardness curves in FIG. 4, suggests the existence of two different mechanisms by which the martensite is softened within the tempering temperature range used here. The first softening mechanism is active below 150° C. and the second mechanism, leading to a sharp drop in hardness, becomes active upon tempering between 200° C. and 250° C. Tempering this armour steels between 200° C. and 250 leads to the coarsening of the metastable transition ε-carbides or η-carbides previously formed below 150° C. and to their transformation into cementite.

Comparison of Mechanical Properties of Experimental Armour Steels

The Charpy impact energies at −40° C., the ultimate tensile strengths and the ratios YS/UTS of the experimental steels are graphically represented in FIGS. 5 to 7, for direct comparison.

Plates of the steels G1A and G1B passed the ballistic testing despite their lower Charpy-V impact energies at −40° C., whereas steel G3** failed.

The higher strength of the steel G2B**, relatively to steel G1A, throughout the entire range of heat treatment parameters applied also is not proportional to the ballistic performance.

However, it appears from FIG. 7 that the steels G1A and G1B that passed the ballistic testing are characterised by relatively lower value of the ratios YS/UTS. FIG. 7 also suggests that lower YS/UTS ratios may be achieved for the steels G2B** and G3** that failed the ballistic testing by increasing their austenitisation temperature to 950° C. and tempering at temperatures lower than 200° C. Such a heat treatment will increase the volume fraction of the retained austenite in the martensite, hence reducing the YS/UTS ratio.

The ratio YS/UTS of the steel G1B (M_(s)=210° C.) remains smaller than 0.68 over the entire range of the austenitisation and tempering temperatures. The Charpy impact energy is slightly higher than in the case of steel G1A but it is lower than the specified 13 Joules within a large range of the heat treatment parameters and becomes higher than 14 Joules for tempering temperatures higher than 100° C. before it again decreases for tempering temperatures higher than 300° C. The values of the UTS higher than the specified 1700 MPa are also found within the same range of tempering temperatures before the decrease below 1600 MPa when the steel G1B is tempered at 300° C. These variations in the Charpy impact energy and UTS are explained hereinafter and are attributed to the effect of the manganese sulphide.

Lower YS/UTS ratios for steel G2A* (M_(s)=255° C.) could be achieved at low normalised tempering temperatures between −1 and 0, or actually lower than 163° C. This limit is lower than the 200° C. found in the case of steel G1A. Steels G1A and G2A* have the same carbon content of 0.39% C but have two different martensite start temperatures of 196° C. and 255° C. respectively, due to the differences in their manganese and chromium contents. The tensile strength of steel G2A* is higher than 1700 MPa throughout the entire range of austenitisation and tempering temperatures used here. This is very high compared to steel G1A and is also high compared to the limit specified for the military and security applications. Steel G2A* also has tensile elongations larger than 11% when tempered at 200° C. This steel has the highest hardness in the as-quenched condition and also after tempering below 200° C. The Vickers hardnesses are above 720 VHN or 640 BHN. Finally it also has a higher YS/UTS ratio than steel G1A.

The volume fraction of retained austenite in the steel G2B** (Ms=271° C.) is lower than the detection limit of about 0.5 volume % that the X-ray diffraction technique could detect.

Some similarities between steel G2B** and steel G2A* may be noted. The YS/UTS ratios are higher than in the case of steel G1A and their tensile strengths are also higher than the specified 1700 MPa throughout the entire range of the austenitisation and tempering temperatures. Their Charpy impact energies at −40° C. are also higher than the specified 13 Joules throughout the entire range of the two heat treatment parameters. Some resemblances are then expected between the microstructures of these two steels that differ from the steel G1A. Their martensite start temperatures are both above 250° C. and no retained austenite (i.e. <0.5%) was detected by X-ray diffraction. The ultimate tensile strength of steel G2B** is slightly lower than for steel G2A* but remains higher than 1700 MPa when the normalised tempering temperature does not exceed the normalised value of 1 or actually 300° C.

The martensite start temperature of steel G3** (M_(s)=309° C.) is higher than the martensite start temperatures of the other four steels. It may be observed that steel G3** represents the highest values of the YS/UTS ratio of all of the five steels considered up to here. The YS/UTS ratio of steel G3** in the untempered condition is in the same range than that of the tempered steels G1A and G1B. The relatively high values of this ratio for steel G3** may be due to an auto-tempering effect during the quenching of this steel, in view of its relatively high M_(s) temperature. The Charpy impact energy of steel G3** is higher than 14 Joules and it remains high throughout the entire range of heat treatment parameters. However the ultimate tensile strength drops to levels lower than 1700 MPa when the tempering temperature is higher than 200° C., while the ratio YS/UTS is higher than 0.68.

General Observations on the Mechanical Properties of Steels G1A Through to G3**.

The five experimental armour plate steels considered here may be classified following their martensite start temperatures, into three groups as shown in Table 6. The first group comprising steels G1A and G1B have martensite start temperatures lower than 210° C. The second group comprises steels G2A* and G2B** and have martensite start temperatures near to 250° C. and the third group comprises steel G3** which has a martensite start temperature near to 300° C.

It may be observed, in FIG. 7, that high martensite start temperatures lead to high values of the YS/UTS ratio in the quenched as well as in the tempered conditions. The YS/UTS ratio increases with an increase in the tempering temperature. The high values of this ratio with high martensite start temperatures, is probably a consequence of auto-tempering during quenching. However, this ratio decreases with an increase in the austenitisation temperature which leads to grain growth and an increase in the volume fraction of retained austenite because of the lower martensite start temperature due to a higher dissolution of carbides. It appears, therefore, that the volume fraction of retained austenite in these armour steels becomes the main factor determining the YS/UTS ratio. The YS/UTS ratio is low with a higher volume fraction of retained austenite and with low tempering temperatures. At higher tempering temperatures both retained austenite and martensite decompose with formation of cementite and ferrite.

TABLE 6 Groups of armour steels classified according to the martensite start temperatures Armour Martensite start Group steel temperature YS/UTS 1 G1A 196° C. <0.65 G1B 210° C. 2 G2A* 255° C. 0.65 to G2B** 271° C. 0.70 3 G3** 309° C. >0.70

Armour steels L300 (M_(s)=285° C.), L500 (M_(s)=253° C.), A66 (M_(s)=241° C.) and MR300 (M_(s)=265° C.) are currently being used in military and security applications within South Africa. The specifications for these armour steels are stated in terms of the yield strength that should be higher than 1500 MPa and the tensile strength, which should be higher than 1700 MPa. These two strength limits will lead to values of the YS/UTS ratio close to 0.88 and will lead to the occurrence of localised yielding during impact. Experience within the industry has shown that steel A66 has better ballistic performance than the other three for plate thicknesses between 8.5 mm up to 30 mm. According to the proposed categories according to M_(s) temperatures, steels L300, L500 and MR300 may be classified into the second group, whereas the steel A66 belongs to the transition between the group 1 and the group 2 of armour steels as previously defined from their martensite start temperatures.

The armour steels in group 3 (M_(s)˜300° C.) have intermediate tensile strengths between those in the first and the second groups. The same observation is valid for their hardnesses. Hence the second group of armour steels is currently produced for military applications based on the design philosophy that would link the expected ballistic performance to the hardness and the tensile strength of these steels, which does not appear to be an entirely valid design philosophy.

The inventors believe that the hardness of armour steel is not an important determinant. Rather the tensile strength is considered to be important as it compares well to the true fracture strength during high-velocity impact. In the present study the YS/UTS ratio is considered in predicting the ballistic performance of the armour steels. These three modes of predicting the ballistic performances using (i) the hardness of the plates, (ii) their ballistic performance index BPI, or (iii) their YS/UTS ratio are compared hereinafter.

Prediction of Ballistic Performances Based on Measured Mechanical Properties

The Ballistic Performance Index (BPI).

The Ballistic Performance Index BPI was introduced by Srivathsa and Ramakrishnan (B. Srivathsa and N. Ramakrishnan, Ballistic performance maps for thick metallic armour, Journal of Materials Processing Technology 96 (1999) 81-91 and B. Srivathsa and N. Ramakrishnan, A ballistic performance index for thick metallic armour, Computer Simulation Modelling in Engineering, 3 (1998), pp. 33-40).

The BPIs of the steels G1A, G1B, G2A* and G2A** calculated according to the above model are shown in Table 7. For this calculation a muzzle velocity of 940 m/s was considered. The average Young's modulus of the steels was assumed to be 200 GPa and the density 7800 kg/m³. The reductions in area used in the BPI and measured by tensile testing, were respectively 6%, 11%, 20% and 8% for these four steels.

TABLE 7 The ballistic Performance Index of the experimental steels G2A* G1A G1B (Failed G2B** (Passed) (Passed) 3/5) (Failed) BPI 3.7 3.9 4.6 4.5 YS [MPa] 880 1100 1500 1500 UTS [MPa] 1780 1897 2200 2000 YS/UTS 0.50 0.58 0.68 0.75

It is inferred from Table 7 that the BPIs of these steels are relatively close to each other but with a tendency to predict a higher ballistic performance for those steels with a higher strength, which contradicts the experimental observation of the ballistic test. The BPI is calculated using mechanical properties measured at lower strain rates and at room temperature. It does not take into account the temperature rise and its dynamic effect upon ballistic impact on the microstructure through phase transformations and transitions in martensitic steels, which, in turn, affect the localised mechanical properties. This can explain the apparent contradiction between the BPI and the actual ballistic performance for these martensitic steels.

The formula for the BPI, however, has the positive value of taking into account the effect of the ductility of the steel on its ballistic performance. It includes the tendency for localised yielding in steels with a high ductility that leads to poor ballistic performance. The BPI formula also demonstrates the decrease in ballistic performance when the velocity of the fired rounds increases. In the case of these steels, the BPI is multiplied by a factor of 3 to 4 when the velocity of the round is reduced from 940 to 400 m/s. But the BPI still predicts a higher ballistic performance for steels that have higher strengths (at least in the case of martensitic steels), which is not necessarily reflected in their actual ballistic performance.

The assessment of performance by ballistic testing remains indispensable and confirms the current observation that a clear relationship between the mechanical properties and the ballistic performance is still lacking. The Ballistic Performance Index should then possibly be considered as a qualitative indication of ballistic performance and may be used for the comparative selection between different armour materials only when their BPIs are different by more than a given margin or a ratio yet to be determined.

The Specifications for South Africa

The current specifications for military and security applications of armour steels in South Africa are:

-   -   the hardness is the main factor determining the ballistic         performance and should be higher than 600 BHN, that is         equivalent to 640 VHN;     -   the transverse Charpy impact energy of the full size specimen         should be higher than 13 Joules at −40° C.;     -   the yield strength of the steel should be higher than 1700 MPa;     -   the ultimate tensile strength should be higher than 2000 MPa;         and     -   the minimum elongation on a 50 mm gauge length is fixed at 6%.

According to this specification the prediction of the ballistic performance was favourable for the steel G2A* only, a steel that resisted two fired rounds but the other three perforated the plate. On the other hand the plates of the steels G1A and G1B passed the ballistic test despite their lower hardnesses and tensile properties than specified. The steel G2B** satisfied all the requirements of this specification except the hardness, and it failed the ballistic test. One should conclude then that the high yield strengths, the high tensile strengths, the high elongations and the high impact energies of the steels G2A* and G2B** did not play a decisive role in resisting high velocity impacts.

The Ratio YS/UTS

The ratios YS/UTS of the experimental tempered martensitic steels are presented in Tables 2 and 7 from which it is observed that the armour steel whose ratios of YS/UTS are lower, had good ballistic performance in the experimental conditions. However, the higher UTS, higher YS or Charpy impact energy, considered separately, appear not to be accurate determinants of the ballistic performance for martensitic steels.

It may be observed that steel G3** (group 3) has the highest impact energy throughout the entire range of the austenitisation and tempering temperatures, whereas steel G1A and steel G1B (group 1) have the lowest impact energy. Steel G2A* (group 2) has a fairly intermediate level of Charpy impact energy. It also appears that the Charpy impact energy of the sub-sized specimens of the armour steels measured at −40° C., increases when the martensite start temperature of the armour steel is higher, which is not the case for the ballistic performance.

Effect of Tempering Temperature and Inclusions on the Mechanical Properties and the Ballistic Performance.

The fracture surfaces after the tensile tests at room temperature and the Charpy impact tests at −40° C. of the three groups of armour steels, as classified according to their martensite start temperatures, were compared by SEM. The effect of silicon, chromium and manganese contents in their resistance to low-temperature tempering, was also analysed. The shear lips near the standard notch of the Charpy specimen as well as the area situated within the fracture surface at a position below the notch and near to the area of contact with the striking edge of the pendulum were analysed by SEM.

It was observed that tempering at 200° C. improves the toughness of these armour steels and increases their Charpy impact energy to above 12 Joules. At the same time the effect of manganese sulphide particles became significant. The fracture face of these Charpy specimens became ductile with small dimples formed near the notch as well as near the impact area away from the notch. The size of the cavities around the manganese sulphide particles became larger as illustrated in FIGS. 8 and 9. The decrease in the Charpy impact energy of the specimens upon tempering above 300° C. may then be partially attributed to the detrimental effect of the decohesion around the manganese sulphide particles in a relatively soft martensite when the tempering temperature exceeds 200° C.

Tempering produces carbides and removes the carbon from solid solution in the martensite and thus lowers the hardness and increases the toughness. However the detrimental effect of the manganese sulphide particles plays a role in the fracture mechanism of these steels in Charpy impact and tensile tests, and imposes a limit to their increase in impact energy with fracture cavities of up to 7 μm that are being formed. The shape of the manganese sulphide particles has a strong effect on the stress concentration effect during Charpy impact and tensile testing. The shape is important but of equal importance here is the very low adhesion between the ferrite matrix and MnS particles. Specifically, upon tempering above 250° C. the softening of the martensite promotes decohesion around the elongated MnS particles. Cavities of diameters larger than 16 μm were formed upon shearing of the areas around the MnS particles. The Charpy impact energy of these armour steels becomes lower once again upon tempering at 400° C.

Both the cementite and the MnS particles are, therefore, prejudicial to the resistance against “lower” strain rate impact loading despite the presence of the soft ferrite. Although not wishing to be bound by theory, this phenomenon can be explained by the occurrence of localised high stresses around the elongated inclusions of the MnS particles. These high stresses are favourable for the nucleation of voids and their coalescence into cavities that consequently lead to a decrease of the nominal ultimate tensile strength of the armour steel upon tempering. The other reason for this decrease of the ultimate strength is the decomposition of the martensite itself and the formation of coarse cementite.

The sub-sized Charpy specimens of steel G2A* (0.009% S, 0.65% Mn) that have martensite start temperatures near to 250° C., show the same brittle behaviour in the untempered condition as was the case with the steels G1A and G1B but with a slightly higher impact energy. Besides the mentioned reasons of the brittle behaviour in the untempered condition, other inclusions such as the calcium-aluminium compounds inherited from the casting process also act as stress raisers and, therefore, may have acted as crack initiators in the hard untempered martensite during the tensile test. For steel G2A* also the effect of the MnS inclusions become observable and large cavities are formed around this type of inclusion that weakened the armour steel when the tensile or the Charpy specimens are tempered at temperatures above 200° C.

The detrimental effect of the MnS particles in the tempered martensite was not observed to have the same importance with the high strain rate ballistic testing. The matrix remains coherent to the inclusions and no cavities or dimples were perceptible around the interfaces as illustrated in FIGS. 10 and 11. This fact may explain the high ballistic performance of steels with relatively lower tensile strengths and Charpy impact energy values. It is clear that the effect of the inclusions on the fracture mechanism depends on the strain rate. At higher strain rates the time is too short for dimples and cavities to initiate and grow with strain around the inclusions or other discontinuities. Rather, grain boundary fracture, with some grains being pulled out of the matrix, was observed to be the failure mode under ballistic impact.

It is inferred from FIGS. 8 to 11 that the fracture modes at “lower” strain rate and at “higher” strain rate (ballistic impact) are different. This may explain the inadequacies between mechanical properties measured by means of tensile or Charpy impact tests and the ballistic performance.

Study 1 Conclusions The brittle fracture of steels G1A and G1B indicates that they cannot be used in the untempered condition because of the risk of spalling if impacted by high velocity projectiles. The tempering treatment at temperatures ranging between 150° C. and 250° C. improves the ductility of the armour steels of Group 1 and 2 at room temperature and at −40° C. SEM also shows that the tempering treatment enhances the negative effect of the MnS particles at “lower” strain rates as in Charpy testing and tensile test. The notch in Charpy testing enhances the brittle behaviour and intergranular fracture of the untempered armour steels in Group 1. All potential stress raisers should therefore be avoided in the manufacture of armour plates.

Inclusions had pronounced negative effects on both the strength and the toughness at “lower” strain rates during tensile and Charpy impact tests than during the high strain rate ballistic impact. The softening of the martensitic matrix by tempering at temperatures above 250° C. enhances the decohesion of the matrix at the inclusion interfaces.

The effect of the inclusions on the fracture mechanism depends on the strain rate. Moreover, phase transformations were observed in the ballistic impacted regions that were not present in tensile or Charpy specimens.

Fracture mechanisms of martensitic steels change drastically between the “lower” strain rate testing and ballistic impact testing. This may explain the lack of correlation between mechanical properties and ballistic performance noticed in the literature. The yield strength, the ultimate tensile strength and the elongation measured at room temperature using “slower” strain rates, and the Charpy impact energy measured at −40° C. are, therefore, not appropriate for the prediction of ballistic performance.

Lower values of the YS/UTS ratio that indicate enhanced resistance to localised yielding, hence higher ballistic performance, are generally obtained with austenitising and tempering temperatures that differ from those necessary for the achievement of high strengths. Strength based design specifications for ballistic performance, are, therefore, not necessarily appropriate on their own and microstructural aspects need to be introduced to improve the prediction of ballistic performance in tempered martensitic armour plate steels.

STUDY 2

In this study an alternative design methodology for tempered martensitic armour steels is proposed which is based on the effect of retained austenite on the ratio of the yield to ultimate tensile strength (YS/UTS), the microstructure of the tempered martensite and its martensite start temperature M_(s). This approach was developed using 6 mm thick armour plates and later was successfully applied to the design of eight experimental armour steels with plate thicknesses ranging from 4.5 to 5.2 mm and tested by a standard 5.56 mm rounds ballistic test.

The effect of the hardness on the ballistic performance of armour steels is not uniquely defined. It depends on the strain rate (function of the striking velocity) and the thickness of the plate. It appears that in thicker plates and under low strain rates the ballistic affected region is more localised than in thinner plates. Therefore, the ability to resist ballistic perforation depends on the hardness in the first case, whereas the ability to deform plastically in a large volume around the impact region becomes the determinant in the case of thinner plate's performance under high strain rates.

The inventors previously studied the thermal effect and the subsequent phase transformations and transitions that occur when high velocity rounds impact on martensitic steel armour plates and which are not accounted for in the BPI model mentioned hereinbefore. Transmission electron microscopy of the ballistic impact regions in eight armour steels suggested that the thermal effect that accompanies the impact and the subsequent phase transformations, might absorb a significant part of the kinetic energy of fired rounds and consequently improve the resistance to ballistic perforation. It would then become necessary to account for these phenomena in the models for an improved prediction of the ballistic performance, allowing the design of higher performance martensitic armour steels.

The design of advanced performance martensitic armour steels from microstructural considerations has the benefit of overcoming the lack of correlation between high strength and high ballistic performance as observed by the inventors.

Materials and Experiments

Chemical Composition and Manufacturing

Five experimental armour steels, namely the steels G1A through to G3 referred to in Study 1 were subjected to standard ballistic testing and their performance compared to those of three currently produced and used armour steels, here named A66, M38 and RL5. Their chemical compositions, casting details, hot rolling processes, heat treatments and the specifications of the ballistic testing applied to these steels are described in Study 1.

Results of the Ballistic Testing

The Ballistic Parameter BP

The results on the ballistic performance of these eight martensitic steels are presented in Table 1 in Study 1.

An alternative Ballistic Parameter (BP) is proposed to account for the microstructural and the plate thickness effects in predicting the ballistic performance, based on earlier experimental results obtained by the inventors on a further 13 martensitic armour steels. The BP is defined as follows:

$\begin{matrix} {{BP} = \frac{{RA}(\%)}{\exp (\delta)}} & (10) \end{matrix}$

where RA is the volume fraction of the retained austenite in the martensitic microstructure and δ is the thickness of the plate in millimetres. The choice of this expression for the BP parameter is based on the proportional lowering of the yield to ultimate tensile strength ratio (YS/UTS) by the presence of retained austenite in the microstructure and on the increase of the effective penetrating mass when the thickness of the plate increases because of the direct transmission of the linear momentum to the cylinder of material ahead of the fired round within the plate.

The Ballistic Parameters of the experimental plates of interest are compared in Table 8.

TABLE 8 Ballistic parameter, YS/UTS and ballistic performance Name of the RA steel [% vol] δ BP YS/UTS G1A 6.0 ± 0.5 6.0 ± 0.2 0.015 0.50 G1B 4.0 ± 0.5 6.0 ± 0.2 0.010 0.58 G2A* 0.6 ± 0.5 6.0 ± 0.2 0.0015 0.68 G2B** 0.6 ± 0.5 6.0 ± 0.2 0.0015 0.75

The 6 mm thick plate of steel G2A* (Group 2) withstood only two of the five rounds fired and, therefore, constituted the demarcating case between those that passed (Group 1) and those that completely failed the ballistic test (Group 3). The Ballistic Parameter BP limit value of 0.010, corresponding to the plate of steel G1B, was then used for subsequent predictions of the ballistic performance of armour plates with thicknesses smaller than 6 mm and was successfully applied to eight further experimental advanced performance steel armour plates with thicknesses ranging from 4.7 mm to 5.2 mm, whilst keeping the ballistic test conditions unaltered.

The BP can be considered as an attempt to find a direct relationship between the microstructure and the ballistic performance instead of an indirect relationship via the mechanical properties. It was observed that plates with BP higher than a value of 0.010 passed the ballistic test, those with a BP near to 0.006 also passed the ballistic test, but dynamic cracks were found to propagate while those plates with a BP˜0.003 failed the ballistic test. It was also observed that the subsequent plastically deformed area around the impact regions increased as the BP increased within the limits of the experiments.

It may also be observed that steels G1A and G1B had values of the yield to ultimate tensile strength ratio (YS/UTS) close to 0.6. These lower values of this ratio indicate the steel's ability to resist localised yielding; in other words, it indicates the ability of the material to dissipate the absorbed kinetic energy through a larger plastic strain around the impact area. This property increases the volume of the material interacting with the fired round, offering better resistance to perforation. The elongation during uniaxial tensile testing indicates the tendency for localised yielding of the steel when impacted. The inventors believe that the elongation during uniaxial tensile testing should be kept lower than 7%. This observation seems contrary to the current specification that recommends an elongation higher than 6% for martensitic armour steel plates thicker than 12 mm.

Microstructure

Thin foil transmission electron micrographs of the armour steels G1A, G1B, G2A* and G2B** before ballistic testing are compared in FIGS. 12 to 14 after a prior tempering at 180° C. for one hour. This tempering temperature was found to be the most optimum for the above three steels (Groups 1 and 2) for achieving a high ductility at room temperature. Fine elongated strings of carbides were found in twinned martensite plates of steels G1A and G1B that were aligned parallel to the martensite plate interfaces as illustrated in FIG. 12 of these two steels that later passed the ballistic test. On the other hand, coarse carbides had precipitated within the martensite laths and on the lath interfaces of steels G2A* and G2B** and these gave poor ballistic performances.

It is inferred from FIGS. 12 to 14 that a high ballistic performance requires a microstructure consisting of twinned martensite with some retained austenite

without coarse carbides. The precipitation of heavy cementite should be avoided by controlling the chemical composition, i.e. through the silicon content of the steel and the tempering temperature.

Steels G1A and G1B that gave successful ballistic performances after quenching and tempering at 150° C. to 250° C. for one hour, were also tempered at higher temperatures of up to 400° C. to find the upper limit of tempering before the detrimental coarse cementite starts to make its appearance. After tempering at 300° C. for 1 hour the TEM thin foil micrographs in FIGS. 15 and 16, show large strings of coarse cementite that formed along the plate interfaces of the steel G1A with 0.21 wt % Si, while in steel G1B with 1.06 wt % Si, noticeably less of these coarse strings of carbides were formed. Hence it is likely that steel G1B will still have a satisfactory ballistic performance, even after tempering at a relatively higher temperature.

The retardation in the formation of coarse cementite during tempering of steel G1B can be attributed to its higher content of 1.06% silicon, an element that is well known for its effect on delaying the formation of cementite from supersaturated metastable martensite. Tempering at 400° C. is not acceptable even for armour steels containing more than 1 wt % silicon because of the significant softening of the material and the formation of coarse cementite, as shown in FIGS. 17 and 18.

With tempering at 400° C., strings of cementite in the steel G1A (with lower % Si) are formed which coarsened along parallel directions contrary to the dispersed particles of cementite that precipitated in the steel G1B with its higher % Si.

The Martensite Start Temperature of the Steel

M_(s) temperatures of the martensitic armour steels were estimated, with an absolute error of ±15° C., using the regression Formula (II), which was based on 23 measured values using dilatometry for armour steels within the range of chemical compositions (in wt %) of interest here.

M _(s)(° C.)=548-590C-35Mn-18Ni-14Cr-9.5Mo-12Si  (11)

The effect of the austenitisation temperature on the martensite start temperature was also analysed for these armour steels in the range between 800 and 950° C., austenitised for 10 minutes in each case. The martensite start temperatures of these armour steels decrease slightly, typically by about 6 to 10° C. when the austenitisation temperature is increased from 800° C. to 950° C. The increase in the austenitisation temperature has many consequences, i.e. greater dissolution of carbides, more solid solution hardening of the parent austenite as well as austenite grain growth and all of these can modify the martensitic transformation process. For instance, greater dissolution of the carbides changes the chemical composition of the matrix and, hence, the chemical driving force for the transformation of the austenite into martensite. It also increases the solid solution hardening of the parent austenite, which affects the movement of the dislocated transformation front through the harder austenite.

The inventors believe that an improved design scheme for high ballistic performance martensitic armour steels should consider the following with regard to the chemical composition and the tempering treatment of the martensitic armour steel:

a. Carbon is the main alloying element that determines the hardness of the martensite. A hardness higher than 500 VHN may be obtained when the carbon content of the armour steel is above 0.37 wt % C.

b. The silicon content of the steel has a strong effect on the stability of the martensite upon tempering, as it delays the softening of the martensite during tempering at higher temperatures. It also appears to increase the resistance to dynamic coarsening of the cementite upon ballistic impact.

Softening of these martensitic steels upon tempering gave the following threshold temperatures (Table 9) where the hardness started to drop measurably:

TABLE 9 Silicon content and the temperature of softening of the martensite in five experimental armour steels. Tempering temperature of measurable Steel wt % Si M_(s) softening G1A 0.21 196° C. 180° C. G3 0.40 309° C. 200° C. G2B 0.43 271° C. 200° C. G2A 0.8 255° C. 250° C. G1B 1.06 212° C. 300° C.

Silicon increases the stability of the martensite by reducing the chemical activity of carbon and, therefore, becomes effective in delaying the decomposition of the martensite in steels in the range between 0.5 wt % to 1.0% C. However, from Table 9 it appears that the stability of the martensite upon tempering does not correlate uniquely with the M_(s) of the steel.

c. The martensite start temperature of these armour steels may be approximated with acceptable accuracy using the experimentally determined regression formula (II).

d. Higher volume fractions of retained austenite in twinned plate martensite lead to lower values of the YS/UTS and to an improved ballistic performance.

A microstructure-based design procedure for advanced performance tempered low-carbon martensitic armour steels, particularly of plates with a thickness of 8.5 mm or less (Group 1) and for high strength components (Group 2) should consist of the following five steps:

Step 1: the Chemical Composition

Select the chemical composition (in wt %) of the steel within the following range:

0.37-0.43% C, 0.5-2.0% Mn, 0.4-1.2% Si, 0.8-1.5% Cr, 0.5-0.6% Mo, 1.8-4.0% Ni with the final chemical composition chosen such that the martensite start temperature corresponds to the following classification and probable areas of application:

TABLE 10 Classification of martensitic armour steels Thickness of armour M_(s) plates Application Group1 <210° C.  4.5-6 mm High ballistic performance Group 2 210° C. < 8.5-20 mm High strength and M_(s) < 260° C. medium ballistic performance Group 3 >260° C.   30 mm High toughness and poor ballistic performance

The M_(s) temperature of the steel may initially be approximated using formula (11), during the design process but later a measured value should be obtained by dilatometry analysis.

Step 2: the Heat Treatment

The austenitisation temperature and the tempering temperature will determine the final properties of the martensitic steels. The austenitisation conditions have an effect on the M_(s) temperature, hence on the thermodynamics and kinetics of the martensite transformation, and on the microstructure.

TABLE 11 Optimum heat treatment and microstructures High performance armour High strength plates components Category Group 1 Groups 2 Austenitisation 870 to 950° C. for 20 to 60 800 to 860° C. for 20 to minutes 60 minutes Quenching Water at room temperature Water or oil medium Tempering 150 to 250° C. for 20 250 to 300° C. for 20 to minutes at least 60 minutes Retained 1-7% volume fraction Not necessary austenite YS/UTS <0.6 >0.75 Microstructure Twinned plate martensite Butterfly and lath with nodular retained martensite without austenite retained austenite Detrimental coarse cementite Elongated manganese particles sulphide, other inclusions and cementite when the tempering temperature increases Strain rate Very high low

Step 3: Prediction of the Ballistic Performance

Perform a phase analysis by XRD, determine the volume fraction of retained austenite and calculate the BP of the plate. The minimum required BP should be equal to 0.010.

Step 4: Assessment of the Performance

Perform a standard ballistic test to confirm the performance prediction on armour plate having BP values higher than 0.010.

Step 5: Microstructure Analysis

Perform TEM and XRD analysis to confirm the dependence of the ballistic performance on the microstructure and phases present.

The proposed design scheme was successfully applied to the design of a further eight experimental martensitic armour steels with plate thicknesses ranging between 4.7 and 5.2 mm and BP values ranging between 0.006 and 0.055.

Study 2 Conclusions

It is possible to predict the ballistic performance of martensitic armour steels by considering the microstructure, morphology and the phases inherited from the combination of chemical composition and heat treatment. The Ballistic Parameter BP, which includes the volume fraction of the retained austenite in these steels and the M_(s) temperature, which determines its morphology, can be considered as criteria for the classification of martensitic armour steels in three application groups.

The microstructure-based design scheme has the advantage of addressing a direct relationship between the microstructure and ballistic performance instead of an uncertain and disproven indirect relationship through the mechanical properties.

For a given chemical composition, it is possible to design for a high ballistic performance or for high strength depending on the heat treatment parameters; a lower YS/UTS ratio indicates a higher resistance to localised yielding upon impact, hence an improved resistance to ballistic perforation for a given composition of the martensitic steel. 

1. A low-carbon martensitic armour steel comprising at least Fe, C, Si and Ni which has a ratio of yield strength to ultimate tensile strength of less than 0.7 and which includes retained austenite at a volume fraction of at least 1%.
 2. The armour steel as claimed in claim 1, which has a ratio of yield strength to ultimate tensile strength of less than or equal to 0.65.
 3. The armour steel as claimed in claim 2, which has a ratio of yield strength to ultimate tensile strength of less than or equal to 0.6.
 4. The armour steel as claimed in claim 1, which comprises 0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni, with the balance being mostly Fe.
 5. The armour steel as claimed in claim 1, which comprises also Mn, Cr and Mo.
 6. The armour steel as claimed in claim 5, which comprises 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight % Mo.
 7. The armour steel as claimed in claim 1, which has a martensite start temperature of less than 210° C.
 8. The armour steel as claimed in claim 7, which has a plate thickness of from 4.5 mm to 8 mm.
 9. The armour steel as claimed in claim 8, which has a plate thickness of from 4.5 mm to 6 mm.
 10. The armour steel as claimed in claim 1, which has a volume fraction of retained austenite of less than 7%.
 11. The armour steel as claimed in claim 1, which has a micro-structure in which the martensite is predominantly present as twinned plate martensite and not lath martensite.
 12. The armour steel as claimed in claim 11, in which the twinned plate martensite is combined with retained austenite in the same micro-structure.
 13. The armour steel as claimed in claim 1, in which any cementite is predominantly present as dispersed particles and not as coarse strings.
 14. The armour steel as claimed in claim 1, which is in the form of a plate with a thickness δmm, and which has a Ballistic Parameter BP of at least 0.01 where BP=volume fraction of retained austenite/exp(δ).
 15. A method of preparing a low-carbon martensitic armour steel, the method including subjecting a steel which comprises at least Fe, C, Si and Ni and which has a martensite start temperature of less than 210° C. to an austenisation heat treatment step at a temperature of at least 800° C.; quenching the steel; and subjecting the steel to a tempering step at a temperature of less than 300° C.
 16. The method as claimed in claim 15, in which the austenisation heat treatment step is at a temperature of between 870° C. and 950° C.
 17. The method as claimed in claim 15, in which the steel is subjected to the austenisation heat treatment step for a period of between 20 minutes and 60 minutes.
 18. The method as claimed in claim 15, in which the tempering step is at a temperature of between 150° C. and 250° C.
 19. The method as claimed in claim 15, in which the steel is subjected to the tempering step for a period of between 20 minutes and 60 minutes.
 20. The method as claimed in claim 15, in which the steel is in the form of a plate with a thickness of from 4.5 mm to 8 mm.
 21. The method as claimed in claim 15, in which the steel comprises 0.37-0.43 weight % C, 0.4-1.2 weight % Si, and 1.8-4.0 weight % Ni, the balance being mostly Fe.
 22. The method as claimed in claim 15, in which the steel comprises also Mn, Cr and Mo.
 23. The method as claimed in claim 22, in which the steel comprises 0.5-2.0 weight % Mn, 0.8-1.5 weight % Cr, and 0.5-0.6 weight % Mo.
 24. A low-carbon martensitic armour steel produced by the method as claimed in claim
 1. 